Water-free titania-bronze thin films with superfast lithium ion transport

ABSTRACT

A multilayered structure including a substrate and a layer of calcium-doped bronze is disclosed. A multilayered structure including a substrate, a layer of calcium-doped bronze, and a layer of pure bronze is also disclosed. A method for fabricating a multilayer structure including a substrate and a layer of calcium-doped bronze is also disclosed.

CROSS-REFERENCE TO RELATED APPLICATIONS

This application claims the benefit of U.S. Provisional Application No. 62/048,115, filed on Sep. 9, 2014. The entire disclosure of the above application is incorporated herein by reference in its entirety.

BACKGROUND

This section provides background information related to the present disclosure which is not necessarily prior art.

Energy storage materials with high capacity and rapid charge/discharge rate are of great interest in lithium ion batteries (LIBs), especially for expanding the application to high power systems such as electric vehicles. The bronze polymorph of titanium dioxide (also referred to as TiO₂—B, TiO₂(b), TiO₂(bronze), titania bronze, “bronze,” etc.) is an excellent candidate due to its open structure and fast lithium ion transport via a pseudocapacitive Faradaic process leveraging ultrahigh discharge rates comparable to those of supercapacitors while maintaining the advantage of storing energy in the bulk. However, existing forms of powder/slurry prepared by conventional hydrothermal methods pose certain challenges, including limited purity, a randomized crystal orientation and the unavoidable presence of lattice water in its structure.

SUMMARY

This section provides a general summary of the disclosure, and is not a comprehensive disclosure of its full scope or all of its features.

The present technology provides for a multilayered structure. The multilayered structure has a substrate and a layer of calcium-doped bronze (Ca:TiO₂—B) in direct contact with the substrate. In certain embodiments, the substrate may be a perovskite material, such as strontium titanate (SrTiO₃). In other embodiments, the substrate may be a non-perovskite material, such as silicon (Si) that is modified with a layer of a perovskite material. In some embodiments, the multilayered structure further comprises a layer of pure bronze (TiO₂—B) deposited on the layer of Ca:TiO₂—B.

In other aspects, the present technology also provides for a multilayered structure that includes a substrate comprising a perovskite material, a layer of calcium-doped bronze (Ca:TiO₂—B) in direct contact with the substrate, and a layer comprising titania-bronze (TiO₂—B) in direct contact with the layer of Ca:TiO₂—B. The layer of Ca:TiO₂—B is positioned between the substrate and the layer of TiO₂—B.

Additionally, the present technology provides for a method for manufacturing a multilayered structure. The method includes depositing a layer of calcium-doped bronze (Ca:TiO₂—B) onto the substrate by pulsed laser deposition (PLD) of calcium titanium oxide (CaTi₄O₉) target onto a substrate. In certain aspects, the PLD comprises laser ablating the CaTi₄O₉ target to generate the Ca:TiO₂—B. In various embodiments, the substrate comprises a perovskite material selected from a group consisting of: SrTiO₃, BaTiO₃, MgSiO₃, CaTiO₃, FeTiO₃, LaMnO₃, PbTiO₃, and mixtures thereof. In other embodiments, the method also includes depositing a layer of pure bronze (TiO₂—B) onto the layer of Ca:TiO₂—B by PLD. The PLD performed for depositing a layer of TiO₂—B may include laser ablating a pure TiO₂ target to generate the TiO₂—B.

Further areas of applicability will become apparent from the description provided herein. The description and specific examples in this summary are intended for purposes of illustration only and are not intended to limit the scope of the present disclosure.

DRAWINGS

The drawings described herein are for illustrative purposes only of selected embodiments and not all possible implementations, and are not intended to limit the scope of the present disclosure.

FIG. 1 shows a perspective view of a first multilayered structure;

FIG. 2 shows a perspective view of a second multilayered structure;

FIG. 3 is a flow chart showing a method for fabricating a multilayer structure according to certain aspects of the present disclosure;

FIG. 4 is a perspective view of a top current collector configuration for electrochemical measurements, wherein films are grown on insulating 10×10 mm² SrTiO₃ substrates;

FIG. 5 shows crystal structures of regular TiO₂—B (top row) and Ca:TiO₂—B (CaTi₅O₁₁, bottom row) projected along [100], [010] and [001] directions, from left to right, which were produced using VESTA;

FIGS. 6A-6D show a characterization of Ca:TiO₂—B and regular TiO₂—B; wherein FIG. 6A is XRD pattern of a Ca:TiO₂—B thin film grown on a (100) SrTiO₃ substrate at 800° C. by pulsed laser deposition (PLD) and the hetero-epitaxial relationship is [100](001)_(Ca:TiO2-B)∥[100](001)_(Sr:TiO3), FIG. 6B is a HAADF STEM image of the Ca:TiO₂—B phase, FIG. 6C is an XRD pattern of a regular TiO₂—B thin film grown on top of a thin Ca:TiO₂—B template layer on a (100) SrTiO₃ substrate under the same PLD conditions, and FIG. 6D is a HAADF STEM image of the dual layer structure;

FIGS. 7A-7B show as-grown morphology of both a Ca:TiO₂ thin film and a TiO₂—B/Ca:TiO₂—B dual layer film on (100) SrTiO₃ substrate, wherein FIG. 7A is a low magnification STEM image of a Ca:TiO₂—B film with about 75 nm thickness, and FIG. 7B is a low magnification STEM image of a TiO₂—B film (about 65 nm) grown on top of a thin Ca:TiO₂—B template layer (about 15 nm), wherein the dash line was drawn to locate the interface;

FIGS. 8A-8E show the structure and electrochemical performance of an inclined Ca:TiO₂—B film grown on a (110) SrTiO₃ substrate, wherein FIG. 8A is a HAADF STEM image of a region near a grain boundary, FIG. 8B is a lower magnification STEM image showing the polycrystalline nature of the film, wherein different grains have channels along different directions exposed at the surface, FIG. 8C is a HAADF STEM image showing a region near the surface after aggressive cycling for more than 60 days, wherein no significant structural degradation, either on the surface or in the film, is observed, FIG. 8D shows charge-discharge voltage profiles of the 5th cycle at each current rate from 1 C to 12000 C (1 C=335 mA g⁻¹), wherein dashed lines depict the profiles at 1 C immediately following the last cycle at 12000 C, and FIG. 8E shows capacity retention at constant 60 C and 80 C rates for 200 cycles in a voltage window of 1-3 V;

FIG. 9 is a graph showing cycling performance of a bare Nb:SrTiO₃ (100) substrate at a 1 C rate for 100 cycles in a voltage window of 1-3 V versus Li metal;

FIGS. 10A-10B show film characterizations, wherein FIG. 10A is a graph showing cyclic voltammograms (CVs) of a (001) TiO₂—B thin film using the top Cu grid current collector at different scan rates, wherein three cycles are shown for each rate, demonstrating good repeatability of the measurement, wherein signature TiO₂—B peaks were observed, and FIG. 10B is a graph showing a specific capacitance calculated by integrating the CVs at different scan rates for both the oxidation and reduction curves;

FIG. 11 is a graph showing a comparison of discharge capacities of a TiO₂—B thin film using two different test configurations, wherein capacity contribution from a Ca:TiO₂—B template layer is subtracted in both configurations, and capacity of the conductive Nb:SrTiO₃ substrate was subtracted for the bottom collector configuration;

FIGS. 12A-12E depict a rate capability comparison and structural change after cycling, wherein FIG. 12A is a graph showing a discharge capacity at the 20th cycle of both TiO₂—B and Ca:TiO₂—B with different orientations at increasing rates, wherein respective film thicknesses are labeled and solid lines are guides for the eyes, the inset of FIG. 12A shows potential profiles at the 5th cycle of (001) TiO₂—B and Ca:TiO₂—B films at C/10 rate (33.5 mA g⁻¹), FIGS. 12B and 12C show before and after cycling structural comparison of a TiO₂—B/Ca:TiO₂—B dual layer film grown on (100) SrTiO₃, wherein a fracture at the interface is clearly seen and such fractures are observed throughout the post-cycling film, and FIGS. 12D and 12E are XRD patterns before and after cycling of both a Ca:TiO₂—B film and a TiO₂—B/Ca:TiO₂—B dual layer film on (100) SrTiO₃;

FIGS. 13A-13B show lattice changes due to Li⁺ intercalation generated by XRD fine scans of (001) bronze films grown on (100) SrTiO₃ substrates, wherein solid curves are fittings of the experimental data, and wherein FIG. 13A is a 006 peak of a Ca:TiO₂—B film with a peak position shift indicating an increase of c and a peak broadening (FWHM: full width at half maximum) were observed, and FIG. 13B shows a TiO₂—B 001 peak of a TiO₂—B/Ca:TiO₂—B dual layer film with a peak position shift indicating a slight decrease of c, and a peak broadening were observed;

FIG. 14 shows fractures throughout a film caused by charging and discharging, wherein a lower magnification STEM image of the same sample as in FIG. 12C shows more fractures in the TiO₂—B/Ca:TiO₂—B dual layer film after cycling, wherein the arrows point to some of the fractures;

FIGS. 15A-15B show a Cu grid surface at a charged state. FIG. 15A is a scanning electron microscopy image showing the surface of the Cu wire on the same Ca:TiO₂—B thin film as in FIGS. 8A-8C, the film being fully charged with Li⁺ at the rate of 1000 C, and removed from the cell at a half cycle to be examined and no obvious Li plating was observed, and FIG. 15B shows an X-ray energy dispersive spectrum taken on the Cu wire, wherein carbon and fluorine are from the residue of LiPF₆ electrolyte;

FIG. 16 shows an image of a multilayered stack having a Pt—TiO₂—B layer and a Ca:TiO₂—B layer;

FIGS. 17A-17C show images of a multilayered stack having a SrTiO₃ substrate, a layer of Ca:TiO₂—B disposed on the substrate, and a layer of Pt—TiO₂—B positioned on the Ca:TiO₂—B layer, where FIGS. 17B and 17C are magnified images of FIG. 17A; and

FIG. 18 is a chart showing comparative electrochemical performance (specific capacity (mAh/cm²) versus current density (mA/cm²)) for comparative electrochemical devices including those having thin layers with platinum (e.g., a first film of Pt—TiO₂—B on (001) SrTiO₃ formed under vacuum conditions and a second film of Pt—TiO₂—B on (001) SrTiO₃ formed with 16 mTorr of oxygen) and thin layers without platinum (e.g., Ca:TiO₂—B on (100) SrTiO₃, TiO₂—B on (100) SrTiO₃, and Ca:TiO₂—B on (110) SrTiO₃).

Corresponding reference numerals indicate corresponding parts throughout the several views of the drawings.

DETAILED DESCRIPTION

Example embodiments will now be described more fully with reference to the accompanying drawings.

The terminology used herein is for the purpose of describing particular example embodiments only and is not intended to be limiting. As used herein, the singular forms “a,” “an,” and “the” may be intended to include the plural forms as well, unless the context clearly indicates otherwise. The terms “comprises,” “comprising,” “including,” and “having,” are inclusive and therefore specify the presence of stated features, elements, compositions, steps, integers, operations, and/or components, but do not preclude the presence or addition of one or more other features, integers, steps, operations, elements, components, and/or groups thereof. Although the open-ended term “comprising,” is to be understood as a non-restrictive term used to describe and claim various embodiments set forth herein, in certain aspects, the term may alternatively be understood to instead be a more limiting and restrictive term, such as “consisting of” or “consisting essentially of.” Thus, for any given embodiment reciting compositions, materials, components, elements, features, integers, operations, and/or process steps, the present disclosure also specifically includes embodiments consisting of, or consisting essentially of, such recited compositions, materials, components, elements, features, integers, operations, and/or process steps. In the case of “consisting of,” the alternative embodiment excludes any additional compositions, materials, components, elements, features, integers, operations, and/or process steps, while in the case of “consisting essentially of,” any additional compositions, materials, components, elements, features, integers, operations, and/or process steps that materially affect the basic and novel characteristics are excluded from such an embodiment, but any compositions, materials, components, elements, features, integers, operations, and/or process steps that do not materially affect the basic and novel characteristics can be included in the embodiment.

Although the terms first, second, third, etc. may be used herein to describe various steps, elements, components, regions, layers and/or sections, these steps, elements, components, regions, layers and/or sections should not be limited by these terms, unless otherwise indicated. These terms may be only used to distinguish one step, element, component, region, layer or section from another step, element, component, region, layer or section. Terms such as “first,” “second,” and other numerical terms when used herein do not imply a sequence or order unless clearly indicated by the context. Thus, a first step, element, component, region, layer or section discussed below could be termed a second step, element, component, region, layer or section without departing from the teachings of the example embodiments. Throughout this disclosure, the numerical values represent approximate measures or limits to ranges to encompass minor deviations from the given values and embodiments having about the value mentioned as well as those having exactly the value mentioned. Other than in the working examples provided at the end of the detailed description, all numerical values of parameters (e.g., of quantities or conditions) in this specification, including the appended claims, are to be understood as being modified in all instances by the term “about” whether or not “about” actually appears before the numerical value. “About” indicates that the stated numerical value allows some slight imprecision (with some approach to exactness in the value; approximately or reasonably close to the value; nearly). If the imprecision provided by “about” is not otherwise understood in the art with this ordinary meaning, then “about” as used herein indicates at least variations that may arise from ordinary methods of measuring and using such parameters. For example, “about” may comprise a variation of less than or equal to 5%, optionally less than or equal to 4%, optionally less than or equal to 3%, optionally less than or equal to 2%, optionally less than or equal to 1%, optionally less than or equal to 0.5%, and in certain aspects, optionally less than or equal to 0.1%.

As referred to herein, ranges are, unless specified otherwise, inclusive of endpoints and include disclosure of all distinct values and further divided ranges within the entire range. Thus, for example, a range of “from A to B” or “from about A to about B” is inclusive of A and of B. Disclosure of values and ranges of values for specific parameters (such as temperatures, molecular weights, weight percentages, etc.) are not exclusive of other values and ranges of values useful herein. It is envisioned that two or more specific exemplified values for a given parameter may define endpoints for a range of values that may be claimed for the parameter. For example, if Parameter X is exemplified herein to have value A and also exemplified to have value Z, it is envisioned that Parameter X may have a range of values from about A to about Z. Similarly, it is envisioned that disclosure of two or more ranges of values for a parameter (whether such ranges are nested, overlapping or distinct) subsume all possible combination of ranges for the value that might be claimed using endpoints of the disclosed ranges. For example, if Parameter X is exemplified herein to have values in the range of 1-10, or 2-9, or 3-8, it is also envisioned that Parameter X may have other ranges of values including 1-9, 1-8, 1-3, 1-2, 2-10, 2-8, 2-3, 3-10, and 3-9.

Titania (TiO₂; also referred to as titanium dioxide) has various polymorphs, which all have different crystallographic properties. The polymorphs include anatase, rutile, brookite, and bronze (TiO₂(B)). As used herein, the bronze polymorph of titania is referred to as “bronze”, “titania bronze” or “TiO₂—B”, wherein “TiO₂” represents titania and “B” represents the “bronze” polymorph. Where TiO₂—B does not contain any dopants or impurities, it may be referred to as “pure TiO₂—B” or “pure bronze”.

In various aspects, the present disclosure provides a calcium-doped titania-bronze material. Titania bronze is a polymorph of titania (TiO₂), which when doped with calcium becomes a novel calcium-doped titania bronze (Ca:TiO₂—B) material. In certain aspects, the present disclosure provides an entirely waterless method of producing phase-pure TiO₂—B in single-crystalline thin films. Such thin film materials have vast applications, including in solar energy conversion, thermoelectrics, photocatalysis, water splitting, and sensors, by way of non-limiting example.

The current technology contemplates a new material, calcium doped bronze (Ca:TiO₂—B; CaTi₅O₁₁). Further, the current technology provides a new waterless process to synthesize hetero-epitaxial crystalline thin films, e.g., comprising Ca:TiO₂—B or TiO₂—B, by using pulsed laser deposition (PLD). In certain aspects, by aligning open channels to out-of-plane directions, extremely high rates of lithium ion transport, up to 12000 C, with extraordinary structural stability can be achieved. As used herein, aligning the open channels to “out-of-plane” directions means aligning the channels in a three dimensional space at an angle from a surface of a substrate or thin film, as opposed to being aligned parallel to a surface of a substrate or thin film. The current technology contemplates forming and utilizing TiO₂—B single crystals. As noted above, materials prepared in accordance with certain variations of the present technology provide new Ca:TiO₂—B and TiO₂—B materials, which are suitable for use in a variety of applications, including as negative electrode materials in LIBs, solar energy conversion, thermoelectrics, photocatalysis, water splitting and sensors.

TiO₂ has been extensively investigated as an anode material for LIBs due to its low cost, minimal environmental impact, structural stability, high theoretical capacity (335 mA h g⁻¹) and inherent safety (a buffer >1.5 V before lithium plating). Fast lithium storage has been demonstrated in anatase, rutile and Li₄Ti₅O₁₂ nanostructures. Although known to have advantages over anatase or rutile, high quality bronze phase titania (TiO₂—B) specimens that demonstrate good electrochemical properties thus far have exclusively been nano-structured powders prepared by hydrothermal methods, as first synthesized in 1980. Being a metastable phase, compounded by the fact that TiO₂-anatase rarely fully reacts and is often used as a precursor in existing synthesis methods, phase pure TiO₂—B has been extremely difficult to obtain, obscuring the interpretation of property testing results. In addition, removal of all H₂O, which could interfere with Li⁺ transport, from the final product is quite difficult, and recent studies have suggested that the presence of H₂O may be needed to keep phase pure TiO₂ from collapsing into anatase upon aggressive heating.

The present technology uses Ca to stabilize the bronze structure without the presence of H₂O, forming a new variant phase Ca:TiO₂—B, and by using Ca:TiO₂—B as a template layer or buffer layer, epitaxial TiO₂—B single-crystalline thin films may be synthesized by PLD, which is a completely waterless process. Significant enhancement in battery performance is achieved by exploiting this epitaxial relationship with the substrate. The ability to accurately control the crystal orientation is especially beneficial to studies focused on surface states, such as in photocatalysis and photovoltaic applications.

With reference to FIG. 1, the present technology provides for a first multilayered structure 10 comprising a substrate 12 and a template layer or buffer layer 14 in direct contact with the substrate 12. The substrate 12 can be composed of any organic or inorganic material commonly used in the art, including crystalline and non-crystalline conductive materials. Non-limiting examples of substrate materials include plastics, polymers, oligomers, metals, silicon (Si), silica (silicon dioxide or silicon oxide,), aluminum oxide, sapphire, germanium, Al₂O₃, MgAl₂O₄, GaP, Ge, InAs MgO, GaAs, silicon carbide, gallium phosphide, gallium nitride, indium phosphide, zinc oxide, aluminum nitride, bismuth germanate, or oxides, nitrides, metals or alloys thereof. In some embodiments, the substrate 12 comprises a perovskite material. In other embodiments, the substrate consists of a perovskite material. As used herein, a “perovskite” material comprises molecules with a perovskite structure having the formula ABX₃, where A and B are cations of different sizes and X is an anion that bonds to both cations. Non-limiting examples of molecules with perovskite structures include SrTiO₃, BaTiO₃, MgSiO₃, CaTiO₃, FeTiO₃, LaMnO₃, PbTiO₃, and combinations thereof. In various embodiments, the substrate 12 is composed of SrTiO₃, (100) SrTiO₃, (110) SrTiO₃. In other embodiments, the substrate 12 is a material comprising a dopant, such as a doped SrTiO₃, doped (100) SrTiO₃, or doped (110) SrTiO₃. Non-limiting examples of suitable dopants include Nb, V, Ru, Rh, Mn, Pd, Ir, and Pt. As a non-limiting example, the substrate can be SrTiO₃ doped with Nb (Nb:SrTiO₃). In embodiments where there substrate 12 does not comprise a perovskite material, such as where the substrate 12 comprises Si, the substrate 12 is modified by with a layer of a perovskite material that serves as a template for the template layer 14. The perovskite material can be any perovskite oxide described herein. As a non-limiting example, a Si substrate can be modified with a layer of SrTiO₃, which is thereby positioned between the substrate 12 and the template layer 14.

In some embodiments, the template layer 14 comprises Ca:TiO₂—B. In other embodiments, the template layer 14 consists of Ca:TiO₂—B. The template layer has a thickness T₁ of greater than or equal to about 1 nm to less than or equal to about 500 nm. The multilayered stack 10 comprising the substrate 12 and template layer 14 can be included in an electrical device without additional films deposited on the template layer 14. However, in some embodiments, not shown in FIG. 1, the multilayered structure 10 further comprises additional layers or films deposited on top of the template layer 14. The additional layers can be electronically active layers or insulating layers.

A second multilayered structure 20 is shown in FIG. 2. The multilayered structure 20 has the same substrate 12 and template layer 14 as the first multilayered structure 10 shown in FIG. 1. However, the second multilayered structure 20 further comprises a TiO₂—B layer 16 comprising TiO₂—B. In some embodiments, the TiO₂—B layer 16 consists of TiO₂—B. In other variations, the TiO₂—B 16 layer may comprise electrically conductive nanoparticles or electrically conductive materials. The TiO₂—B layer 16 is in direct contact with the template layer 14. A near-perfect lattice match between the surface structures of the template layer 14 and the TiO₂—B layer 16 provides for a high stability and efficient Li⁺ transport between the layers 14, 16. As used herein, a “near-perfect lattice match” means the interatomic periodic spacing of the template layer's lattice structure is about the same as the interatomic periodic spacing of the TiO₂—B layer's lattice structure at their interface. The TiO₂—B layer 16 has a thickness T_(B) of from about 1 nm to about 1000 nm. In some embodiments, the TiO₂—B layer 16 has a thickness T_(B) of from about 50 nm to about 200 nm. As in the first multilayered structure 10, the substrate 12 of the second multilayered structure 20 is modified with a layer comprising a perovskite material, such as a perovskite oxide, when the substrate 12 does not comprise a perovskite structure. In some embodiments, not shown in FIG. 2, the multilayered structure 20 further comprises additional layers or films deposited on top of the TiO₂—B layer 16. The additional layers can be electronically active layers or insulating layers.

In various embodiments, such as for LIBs, the first and second multilayered structures 10, 20 may be incorporated into an electrochemical cell assembly. In such a variation, the device has first and second multilayered structures 10, 20 that further comprise a wire grid, a separator layer or membrane, a counter electrode, a nonaqueous electrolyte, and a casing, such as a stainless steel cell casing. The wire grid is composed of any conducting material commonly used in the art, such as, for example, copper, gold or platinum. The separator is composed of any material commonly used in the art for battery separators, such as, for example, nonwoven fibers, polymer films, and naturally occurring substances. Non-limiting examples of nonwoven fibers include cotton, nylon, polyesters, and glass; non-limiting examples of polymer films include polyethylene, polypropylene, poly(tetrafluoroethylene), and polyvinyl chloride; and non-limiting examples of naturally occurring substances include rubber, asbestos, and wood. In various embodiments, the multilayered structures 10, 20 are incorporated in a half-cell, wherein the counter electrode is an anode and the Ca:TiO₂—B and/or the TiO₂—B are the cathode. In a half cell, the anode is a metal layer, such as, for example, Li metal, Na metal, K metal, or Mg metal, as non-limiting examples. In other embodiments, the multilayered structures 10, 20 are incorporated in a full-cell, wherein the Ca:TiO₂—B and/or the TiO₂—B are an anode, and the counter electrode is a cathode comprising a material with a higher potential than the Ca:TiO₂—B and/or the TiO₂—B anode, such as, for example, LiCoO₂, LiNi_(0.5)Mn_(1.5)O₄, or LiFePO₄. The non-aqueous electrolyte may comprise a lithium salt in an organic solvent, such as ethylene carbonate, dimethyl carbonate, diethyl carbonate, and combinations thereof as non-limiting examples. Non-limiting examples of the lithium salt include LiPF₆, LiBF₄ and LiClO₄.

With further reference to FIGS. 1 and 2, in certain other embodiments, a plurality of conductive nanoparticles may be embedded in or otherwise distributed throughout the template layer 14 and/or in the TiO₂—B layer 16. Embedding or otherwise distributing conductive nanoparticles, for example, in the TiO₂—B layer 16 can significantly improve the electrical conductivity of the TiO₂—B layer, which results in enhanced performance when used in a battery or other electrochemical device relative to a battery or other electrochemical device lacking conductive nanoparticles. A TiO₂—B layer comprising conductive nanoparticles is also useful for photocatalysis water splitting applications.

In certain variations, the conductive nanoparticles may have a particle size (an average diameter for the plurality of nanoparticles present) of greater than or equal to about 10 nm to less than or equal to about 100 nm. The conductive nanoparticles may be formed of a variety of conductive materials including metallic, semiconducting, ceramic, and/or polymeric nanoscale particles having plurality of shapes. In certain variations, the nanoparticles may comprise conductive metal materials like platinum, gold, silver, copper, aluminum, nickel, iron, cadmium, mercury, lead, molybdenum, iron, and alloys or compounds thereof. Particularly suitable nanoparticles comprise platinum. In other alternative variations, suitable conductive nanoparticles can be exemplified by, but are not limited to, graphene/graphite, carbon (such as carbon nanotubes, like single walled nanotubes (SWNTs) or multi-walled nanotubes (MWNTs)), silicon, seedling metals, CdTe, CdSe, CdS, HgTe, HgSe, HgS, PbTe, PbSe, PbS, MoS₂, FeS₂, FeS, FeSe, WO_(3-x), and other similar materials known to those of skill in the art.

Platinum is a particularly suitable conductive material for use as a nanoparticle for the generation of platinum containing TiO₂—B (Pt—TiO₂—B). For example, Pt—TiO₂—B may be used as a template layer 14, on top of which is positioned a layer of TiO₂—B 16, or the Pt may be embedded in the TiO₂—B layer 16 to generate a Pt—TiO₂—B layer 16 adjacent to a template layer 14 comprising Ca:TiO₂—B.

The nanoparticles may vary in concentration within the layer from a first surface of a conductive nanoparticle-containing TiO₂—B material layer to a second opposing surface of the conductive nanoparticle-containing TiO₂—B material layer. For example, in some embodiments, a concentration gradient is formed, where a relative concentration of conductive nanoparticles increases from the first surface to the second opposing surface of the conductive nanoparticle-embedded TiO₂—B layer. In other embodiments, the relative concentration of conductive nanoparticles in the layer may decrease from the first surface to the second opposing surface of the conductive nanoparticle-embedded TiO₂—B material layer. In yet other embodiments, there is a relatively low concentration of the conductive nanoparticles near the first surface and second opposing surface and a relatively high concentration of the conductive nanoparticles between the first and second and second opposing surface of the conductive nanoparticle-embedded TiO₂—B layer.

With reference to FIG. 3, the present technology also provides a method 30 for producing the multilayered structures 10, 20. As shown in box 32, the method comprises obtaining a substrate. The substrate can be any substrate described above in regard to the current technology. Accordingly, the substrate may comprise a perovskite material or a non-perovskite material modified by a perovskite material. For example, in various embodiments, the substrate comprises silicon (Si), which may be modified with a perovskite material, such as SrTiO₃, or other perovskite oxide, by methods known in the art, such as by molecular-beam epitaxy (MBE). As a non-limiting example, modifying a Si wafer substrate with SrTiO₃ comprises depositing half a monolayer (ML=6.8×10¹⁴ atoms/cm²) of strontium on a clean Si (001) surface at a substrate temperature of about 700° C. The wafer is then cooled to near room temperature where an additional ½ ML of strontium is deposited under ultra-high vacuum (UHV). Oxygen is then introduced and additional strontium is deposited in the presence of the oxygen to form a total of 3 crystalline SrO MLs. On top of the 3 MLs of crystalline SrO, 2 ML of amorphous TiO₂ is deposited in oxygen with the substrate temperature still near room temperature (under 200° C.). The resulting heterostructure is then annealed in UHV at equal to or greater than about 450° C. to equal to or less than about 550° C. to recrystallize a SrTiO₃ layer 2.5 unit cell thick. Further growth of the epitaxial SrTiO₃ layer to a desired film thickness can be achieved on the recrystallized 2.5 unit-cell-thick SrTiO₃ template layer through the repeated co-deposition (Sr+Ti+O₂ molecular beams) of an amorphous SrTiO₃ layer near room temperature, followed by recrystallization in UHV. Such a method is reported by Mi et al., “Atomic structure of the interface between SrTiO₃ thin films and Si (001) substrates,” Appl. Phys. Lett. 93, 101913 (2008), which is incorporated herein by reference in its entirety.

As shown in box 34, the method 30 further comprises depositing a Ca:TiO₂—B template layer onto the substrate. According to the method 30, the Ca:TiO₂—B is grown from a calcium titanium oxide (CaTi₄O₉) target, which may be made by preparing a mixture comprising 50%-90% (wt) TiO₂ powder and 10%-50% (wt) CaO powder or 10%-50% (wt) CaTiO₃ powder, sintering the mixture at a temperature from about 1000-2000° C. to generate a powdered material, and pressing the powdered material into a pellet under from about 5000-15,000 lbs of force to generate the CaTi₄O₉ target. In one embodiment, the Ca:TiO₂—B is grown from a CaTi₄O₉ target made by preparing a mixture comprising 80% (wt) TiO₂ and 20% (wt) CaO powders, sintering the mixture at about 1400° C. to generate a powdered material, and pressing the powdered material into a pellet under 10,000 lbs of force. The Ca:TiO₂—B template layer is deposit by PLD in a vacuum chamber with a base pressure of <10⁻⁷ Torr. Ca:TiO₂—B deposition is performed with an about 200 nm to about 300 nm excimer laser with a pulse duration of from about 10 ns to about 50 ns, a fluence of from about 2 to about 5 J cm⁻², a repetition rate of from about 2 Hz to about 20 Hz, and at a substrate-target distance of from about 5 to about 10 cm. Thin Ca:TiO₂—B films are deposited at from about 500 to about 1000° C. in an oxygen ambient of from about 0.025 to about 0.075 Torr, at a deposition rate of from about 0.005 Å/pulse to about 0.05 Å/pulse. For example, in one embodiment, Ca:TiO₂—B deposition is performed with a 248 nm KrF excimer laser with a pulse duration of 22 ns, a fluence of about 3.4 J cm⁻², a 10 Hz repetition rate, and a substrate-target distance of 6.35 cm, and thin TiO₂—B films are deposited at 800° C. in an oxygen ambient of 0.05 Torr, and at a deposition rate of from about 0.01 Å/pulse to about 0.02 Å/pulse. As shown at 36, the method 30 stops in embodiments where no further layers are desired, or when no TiO₂—B layer is desired.

As shown in step 38, in some embodiments the method 30 also comprises depositing a TiO₂—B layer onto the Ca:TiO₂—B template layer. Depositing the TiO₂—B layer onto the Ca:TiO₂—B template layer is performed by ablating a pure TiO₂ target under the same conditions described above for depositing the Ca:TiO₂—B template layer on the substrate. Although PLD is described above, in various embodiments the template layer and layer of pure bronze may be deposited by other techniques, such as sputtering, atomic layer deposition (ALD), chemical vapor deposition (CVD), or by molecular beam epitaxy (MBE).

Embodiments of the present technology are further illustrated through the following non-limiting examples.

Example 1

Methods for Generating Multilayered Stacks and Half-Cells

A CaTi₄O₉ target used to grow Ca:TiO₂—B thin films is made by mixing 80% TiO₂ and 20% CaO powders, sintering at 1400° C., and pressing into a pellet under 10,000 lbs of force. A vacuum chamber used for PLD has a base pressure <10⁻⁷ Torr. A 248 nm KrF excimer laser with a pulse duration of 22 ns and a fluence of ˜3.4 J cm⁻² is used for the deposition at a 10 Hz repetition rate, and the substrate-target distance is set to 6.35 cm. Thin films are deposited at 800° C. in an oxygen ambient of 0.05 Torr. The deposition rate is 0.01-0.02 Å/pulse. Deposited films have thicknesses of 50-200 nm (typical deposition time of 1-4 hours), which are measured by a Veeco Dektak profilometer and confirmed with TEM images. XRD results are obtained on a Rigaku rotating anode diffractometer using Cu Kα radiation. All STEM images are captured on a JEOL 2100F TEM equipped with a spherical aberration corrector. Devices are also made with a film of pure TiO₂—B grown on the Ca:TiO₂—B film. The pure TiO₂—B film is deposited on the Ca:TiO₂—B film by the same protocol described above for the Ca:TiO₂—B film, but with pure 100% TiO₂ as a target instead of CaTi₄O₉.

Battery half-cells (EL-CELL ECC-STD) are assembled in an argon-filled glove box (Innovative Technology Inert Lab) with O₂ and H₂O levels below 2 and 1 ppm, respectively, and tested at room temperature on a Princeton Applied Research VersaSTAT MC 4-channel system operating in galvanostatic mode using a lithium metal anode, a non-aqueous electrolyte (1M LiPF₆ in ethylene carbonate:dimethyl carbonate 1:1 (v/v), Merck) and a 1.55 mm thick glass fiber separator. In order to investigate the electrochemical performance of thin films, a current collector is needed, either on top or at the bottom of the film. Conductive SrTiO₃ substrates doped with 0.5 at. % Nb (resistivity 0.05 Ωcm) are used as bottom current collectors, and provide similar film quality to those grown on undoped SrTiO₃ substrates. Conversely, for films grown on non-conductive SrTiO₃ substrates, a top current collection geometry is fabricated using a grid of Cu wires with line width of 100 μm and a thickness of 20 nm which are deposited on the film surface in an E-beam evaporator with a Mo mask. On a 10×10 mm² test sample, the grid covers <0.8% of the surface area, which has a negligible influence on the Li⁺ exchange between the film and the electrolyte. Such a configuration is shown schematically in FIG. 4. Although not shown in FIG. 4, some half-cells comprise only the substrate and the Ca:TiO₂—B film, and not the pure TiO₂—B film.

Results and Characterizations

The CaTi₅O₁₁ stable phase is discovered in atomic resolution high-angle annular dark-field (HAADF) scanning transmission electron microscopy (STEM) images, according to which a geometric model is built and first-principles optimization are performed by PW91 functional implemented in VASP. More accurate HSE06 method calculations indicate that CaTi₅O₁₁, or Ca₄Ti₂₀O₄₄ in a unit cell, is an orthorhombic structure with the symmetry of CMCM (63) and lattice constants of a=12.1702 Å, b=3.8013 Å, c=17.9841 Å, α=β=γ=90°. Atom positions are shown in Table 1. The PLD target recipe and the growth conditions, which are detailed above, produce a high quality CaTi₅O₁₁ thin film deposited onto a (100) SrTiO₃ substrate. The crystal structure projected along three crystallographic directions and is compared with the regular TiO₂—B structure in FIG. 5. FIG. 6A shows the θ-2θ X-ray diffraction (XRD) pattern for an epitaxial (001) thin film of the CaTi₅O₁₁ phase. A HAADF STEM image taken along the b direction of the structure is displayed in FIG. 6B, showing an interesting layered, zigzag pattern where every inserted layer of alternating Ti and Ca atoms flips the stacking direction of the next two layers of Ti atoms above it, consistent with the atomic model in FIG. 5. Because it is a variant of the TiO₂—B structure with extra Ca layers and superlattice twinning, the new phase is designated as Ca:TiO₂—B.

TABLE 1 Atom positions in the Ca:TiO₂—B structure # Atom x y z 1 O O1 0.19999 0.93727 0.34267 2 O O2 0.04096 0.93438 0.98907 3 O O3 0.23305 0.93649 0.11333 4 O O4 0.07177 0.93646 0.21254 5 O O5 0.38543 0.93642 0.24343 6 O O6 0.7005 0.43782 0.34213 7 O O7 0.53992 0.43719 0.98906 8 O O8 0.73327 0.43714 0.11332 9 O O9 0.57158 0.43772 0.2121 10 O O10 0.88567 0.43771 0.24335 11 O O11 0.78487 0.9355 0.99189 12 O O12 0.94481 0.93819 0.34546 13 O O13 0.75278 0.93744 0.22097 14 O O14 0.91413 0.93789 0.12209 15 O O15 0.6003 0.93703 0.09083 16 O O16 0.28498 0.43691 0.9919 17 O O17 0.44504 0.43607 0.34505 18 O O18 0.25261 0.43655 0.22121 19 O O19 0.41393 0.43555 0.1219 20 O O20 0.09999 0.43618 0.0913 21 O O21 0.23209 0.93441 0.72099 22 O O22 0.07037 0.93436 0.62223 23 O O23 0.38436 0.93492 0.59074 24 O O24 0.73211 0.43547 0.72103 25 O O25 0.57044 0.43497 0.62189 26 O O26 0.88417 0.43548 0.59087 27 O O27 0.75136 0.93542 0.61322 28 O O28 0.2515 0.43522 0.61321 29 O O29 0.41305 0.43458 0.71227 30 O O30 0.0992 0.43408 0.74361 31 O O31 0.91301 0.93537 0.71234 32 O O32 0.59921 0.93577 0.74336 33 O O33 0.94383 0.9365 0.48925 34 O O34 0.69927 0.43558 0.49164 35 O O35 0.44389 0.43501 0.48876 36 O O36 0.19931 0.93896 0.49221 37 O O37 0.54041 0.4367 0.84527 38 O O38 0.28431 0.43673 0.8422 39 O O39 0.03923 0.93309 0.84518 40 O O40 0.78454 0.93558 0.84229 41 O O41 0.5666 0.93536 0.41677 42 O O42 0.41789 0.93805 0.91639 43 O O43 0.91796 0.43314 0.91749 44 O O44 0.06658 0.43879 0.41808 45 Ti Ti1 0.22043 0.93658 0.24214 46 Ti Ti2 0.06116 0.93641 0.09019 47 Ti Ti3 0.72074 0.43715 0.24147 48 Ti Ti4 0.56112 0.43659 0.09009 49 Ti Ti5 0.76538 0.93697 0.09239 50 Ti Ti6 0.92466 0.93757 0.2442 51 Ti Ti7 0.26493 0.43673 0.09246 52 Ti Ti8 0.42448 0.43642 0.24382 53 Ti Ti9 0.21922 0.93567 0.59265 54 Ti Ti10 0.06009 0.93443 0.74408 55 Ti Ti11 0.71913 0.43538 0.59214 56 Ti Ti12 0.56028 0.43575 0.74408 57 Ti Ti13 0.26434 0.43449 0.7417 58 Ti Ti14 0.42346 0.43479 0.58992 59 Ti Ti15 0.76407 0.93521 0.74178 60 Ti Ti16 0.9231 0.93527 0.59028 61 Ti Ti17 0.58362 0.45134 0.41619 62 Ti Ti18 0.40087 0.42763 0.91679 63 Ti Ti19 0.90097 0.9446 0.91686 64 Ti Ti20 0.08325 0.93465 0.41732 65 Ca Ca1 0.87195 0.43614 0.41755 66 Ca Ca2 0.37179 0.93857 0.41769 67 Ca Ca3 0.61261 0.93498 0.91722 68 Ca Ca4 0.11262 0.43539 0.91721

This Ca:TiO₂—B structure is used alone or as a template layer to grow (001) regular TiO₂—B thin films because direct deposition of pure TiO₂ on SrTiO₃ substrates usually results in the anatase phase, but highly crystalline TiO₂—B forms on the a-b plane of a Ca:TiO₂—B layer by ablating a pure TiO₂ target under the same growth conditions, mainly because of the near-perfect lattice match between the two phases. The morphology of both the Ca:TiO₂—B thin film and the TiO₂—B/Ca:TiO₂—B dual layer film on a (100) SrTiO₃ substrate are shown in FIG. 7. As shown in FIG. 7, both the Ca:TiO₂—B film and the TiO₂—B/Ca:TiO₂—B dual layer film have fairly smooth surfaces on the a-b plane due to the characteristically layered structure of bronze. Crystal defects including grain boundaries, dislocations and stacking faults are identifiable, but minimal. The structure parameters of the TiO₂—B films are in good agreement with values in the literature. A template layer with a thickness of 10 nm is developed, and no degradation of the TiO₂—B crystalline quality is observed. FIGS. 6C and 6D show the θ-2θ XRD pattern and the HAADF image, respectively, of the dual layer structure. The theoretical densities of Ca:TiO₂—B and TiO₂—B are determined to be 3.637 g cm⁻³ and 3.616 g cm⁻³, respectively, approximately 7% lower than that of anatase.

FIG. 5 shows the crystal structures of both regular TiO₂—B and its variant Ca:TiO₂—B phases. Although the extra Ca atoms make the latter more complicated, there is still a strong resemblance between the two structures. The STEM images in FIGS. 6B and 6D are clearly associated with the two atomic structures projected along the [010] direction in the center column of FIG. 5. The geometric model of the novel Ca:TiO₂—B phase is refined by analyzing the HR-STEM images and performing first-principles optimization. Again, the atom positions in the Ca:TiO₂—B structure are shown in Table 1.

With further reference to FIG. 5, by visual inspection, the channels running along the b-axis appear to be most open among the three crystallographic axes in both structures, and are a good candidate for high Li⁺ mobility in the crystal. Channels parallel to various other directions can are also found by manipulating the model. It is worth noting that more rigorous study than simply observing the cross sectional areas of the channel opening is required to determine with relative certainty the actual diffusion path that is most energetically favorable for fast Li⁺ transport, as the results may be counter-intuitive. The ability to fabricate crystalline thin films of the active storage material with well-defined lattice plane on the surface, such as those described herein, is therefore of great value to experimentally determine the preferred Li⁺ pathways.

For both the Ca:TiO₂—B and TiO₂—B structures, Li⁺ access into the crystal is expected to be easier in the a-b plane, i.e., through the well-aligned channels along the a-axis between layers of atoms (FIG. 5, and FIGS. 6B and 6D) as well as the possibly even faster channels along the b-axis, than in the perpendicular direction. Therefore, in some embodiments these open channels are exposed at the film surface to increase the rates of lithium ion transport, which in principle can be achieved by utilizing substrates with a different orientation. FIG. 8B shows the HAADF image of a Ca:TiO₂—B film deposited on a (110) SrTiO₃ substrate under the same growth conditions as above. Instead of being parallel to the surface, the a-b planes are now inclined, with channels along the a- and b-axes reaching the surface. FIG. 8A clearly displays a region near the boundary between two such grains, where one grain (right) has channels parallel to the a-axis running to the surface, and another grain (left) is rotated about the a-b plane normal having channels parallel to the b-axis running to the surface. FIG. 8B shows the polycrystalline nature of the film. However, such a Ca:TiO₂—B film does not serve as a good template layer for the subsequent growth of uniform TiO₂—B on top. Nonetheless, TiO₂—B is grown on the film.

To characterize their electrochemical properties, the thin films are assembled in half-cells with metallic Li as a counter electrode. To calculate specific capacity, the mass of active material is determined from its theoretical density, measured surface area and thickness. The mass loading of active material is about 0.036 mg/cm². Cycled cells are disassembled in the glove box, and post-cycling films are washed in dimethyl carbonate for three times and dried in vacuum overnight before XRD and TEM studies.

While using a conductive Nb:SrTiO₃ substrate as bottom current collector, electrons travel through the entire substrate to the external circuit, and therefore the electrochemical force may drive some Li⁺ into the substrate. Even though SrTiO₃ does not appear to have a high Li⁺ capacity, it is important to rule out the contribution from the substrate for determining the actual capacity of the film. Therefore, a bare Nb:SrTiO₃ substrate is assembled in a half-cell and tested with exactly the same routine and rates as for the thin film samples. The measured capacity of the substrate at each rate is then subtracted from the total to obtain the capacity of the film at that rate. It should be noted that the voltage window of 1-3 V for TiO₂—B film testing is much higher than the possible Li intercalation voltage of SrTiO₃, so the substrate contribution is very low, as seen in FIG. 9, which shows the control test at a 1 C rate.

A similar approach is needed to determine the capacity of the regular TiO₂—B phase. Since the TiO₂—B film is grown on top of a Ca:TiO₂—B template layer, the method for determining its specific capacity is to cycle the Ca:TiO₂—B sample and the TiO₂—B/Ca:TiO₂—B dual layer sample of the same sizes using exactly the same routine and rates, determine the specific capacity of Ca:TiO₂—B at each rate first, calculate the capacity contribution of the Ca:TiO₂—B layer in the dual layer sample from its thickness obtained by STEM, and finally subtract that part from the total capacity.

For thin films grown on insulating SrTiO₃ substrates and using the top current collection configuration, the substrate is not a part of the electrochemical reaction or the circuit, so its contribution to the measured capacity should be minimal and hence is not considered. To test the effectiveness of the experimental setup, cyclic voltammograms (CVs) of the TiO₂—B thin film sample are recorded at scan rates from 0.1 to 1 mV s⁻¹ as shown in FIG. 10A. A pair of redox peaks at 1.54 V and 1.69 V is observed, which represent the signature pseudocapacitive Li⁺ storage behavior of TiO₂—B. The specific capacitance is calculated by integrating the CVs and the calculated capacitance is almost the same from either the oxidation or the reduction curve, regardless of scan rate, which also corresponds well with the specific capacity obtained from galvanostatic cycling, as shown in FIG. 10B. Such results directly prove the validity of the current testing methods.

For the purpose of comparing the effectiveness of these two configurations described above, two TiO₂—B/Ca:TiO₂—B dual layer control samples are grown simultaneously to the same thicknesses on a (100) SrTiO₃ substrate and a (100) 0.5 at. % Nb:SrTiO₃ substrate, respectively. XRD and TEM results confirm that the two films are of about equal quality. The former is processed in a top current collection geometry as in FIG. 4. Both samples, together with a bare Nb:SrTiO₃ substrate, are then assembled in half-cells with metallic Li counter electrodes and tested under 1 C and 10 C rates for 100 cycles. The discharge capacities of TiO₂—B obtained in the two samples are compared in FIG. 11. It is clear that the two test configurations produce similar results, while the capacity values acquired from the top Cu grid method are slightly higher. This is due to the fact that Cu metal provides better current collection efficiency than Nb:SrTiO₃ semiconductor substrate. Such an advantage becomes more significant as the rates increase beyond 10 C. Therefore, results of battery cycling performance reported herein are all acquired in the top Cu grid collector configuration.

For Ca:TiO₂—B, assuming 5 Li⁺ is intercalated per CaTi₅O₁₁ formula unit (making all Ti 3⁺), its theoretical capacity is estimated to be 294 mA h g⁻¹. For simplicity and comparison with TiO₂—B is defined as 1 C=335 mA g⁻¹. Superior charge/discharge rate capability is observed in the Ca:TiO₂—B film grown on (110) SrTiO₃ with open channels extending to the surface. Starting at 1 C, the battery half-cell was charged and discharged between 1 and 3 V for 50 cycles at each of several rates up to an extreme of 12000 C, ending again at 1 C immediately following the last cycle at 12000 C for additional 20 cycles to examine the structural stability. FIG. 8D shows the voltage curves of the 5th cycle at each rate. At 1 C, the film discharges to a specific capacity of 293 mA h g⁻¹, over 99.6% of the theoretical capacity. The capacity reduced to 248 mA h g⁻¹ at 10 C, 61.4 mA h g⁻¹ at 120 C, and 28.8 mA h g⁻¹ at 600 C. At 12000 C the capacity was 11.5 mA h g⁻¹, likely because only a fraction of the film close to the surface and the current collector is actually lithiated. When the rate was lowered back to 1 C, the capacity was immediately restored to 284 mA h g⁻¹, showing outstanding endurance of the material under extreme conditions. The majority of the capacity occurred in the sloped regions of the voltage profiles, while the specific capacity obtained by integrating the cyclic voltammograms was almost the same regardless of the scan rate (FIGS. 10A-10B), expressing the pseudocapacitive Faradaic behavior of Li storage in this material. The battery half-cell was continuously charged and discharged for 200 cycles at 60 C and 80 C, as shown in FIG. 8E, delivering discharge capacities of 155 mA h g⁻¹ and 102 mA h g⁻¹ at the 100th cycle, respectively. The capacity loss and the lower Coulombic efficiency in the first 10 cycles reflect the poor electrical conductivity characteristic of pure Ca:TiO₂—B. From cycle 10 to 200, however, the loss was only 0.1% per cycle, and the Coulombic efficiency close to 1.

Taking advantage of a clearly defined lattice orientation, the presumed preference for Li⁺ transport along certain crystal directions is demonstrated. The rate capability of Ca:TiO₂—B thin films grown on (110) SrTiO₃ with channel openings on the surface is compared with both Ca:TiO₂—B and TiO₂—B grown on (100) SrTiO₃ with channels along a- and b-axes running parallel to the surface. Considering the impaired electron transport in these materials, all three samples are grown to almost the same thickness in order to ensure a fair comparison. Slow cycling tests at a C/10 rate shown in the inset of FIG. 12A are performed on the two (001) films that have in-plane a- and b-channels, where TiO₂—B and Ca:TiO₂—B discharge to 334 (Li_(0.997)TiO₂) and 273 mA h g⁻¹, respectively. As the rate increases (FIG. 12A), TiO₂—B delivers higher capacities than Ca:TiO₂—B at every rate, indicating that Li⁺ transport along the out-of-plane direction is faster in TiO₂—B than in Ca:TiO₂—B, because the difference in theoretical capacity alone is unlikely to account for such discrepancy. On the other hand, the Ca:TiO₂—B film with exposed a- and b-channels exhibits far superior rate capabilities to both of the above, suggesting a better efficiency of inserting and extracting Li⁺. The superiority becomes more and more significant with increasing rates. For example, its capacity at 60 C is even higher than the 10 C capacity of TiO₂—B and the 1 C capacity of the same Ca:TiO₂—B phase in the different crystal orientation. These results suggest that Li⁺ transport into the bulk of the material is indeed much faster in the Ca:TiO₂—B film on (110) SrTiO₃, either because channels within the a-b plane are more favorable for Li⁺ transport than those along the c-axis, or due to certain effects associated with the polycrystalline structure and grain boundaries.

Slow charge/discharge cycling experiments at a C/10 rate are performed on a Ca:TiO₂—B film and a TiO₂—B film (with a thin Ca:TiO₂—B template layer) both grown on (100) SrTiO₃ substrates using the top Cu grid collector as shown in FIG. 4. The voltage profiles are displayed in the inset of FIG. 12A. Both samples exhibit sloped profiles corresponding to a pseudocapacitive process of Li⁺ transport, which is a typical characteristic often observed for TiO₂—B. These results also support the conclusion that both the TiO₂—B and the Ca:TiO₂—B films have good purity without a pronounced amount of other TiO₂ polymorphs, which would otherwise create plateaus in the profiles.

The material's response to the intensive cycling is examined by XRD and transmission electron microscopy (TEM). As shown in FIGS. 12D and 12E, XRD patterns of both the TiO₂—B and Ca:TiO₂—B structures before and after being cycled for over 40 days are essentially unchanged. Atomic resolution TEM analysis confirms that all the bronze structures stay intact without any signs of significant degradation or collapse into other TiO₂ polymorphs. As an example, a post-cycling HAADF image of the inclined Ca:TiO₂—B structure is shown in FIG. 8C. Relatively subtle changes, however, are observed by comparing the images before and after cycling in the horizontally oriented (001) films. FIG. 12B shows a region near the interface between TiO₂—B and Ca:TiO₂—B template in an as-grown dual layer film on (100) SrTiO₃. A straight, inclined anatase “wall” about 3 nm wide stems from the interface and extends to the surface, separating two TiO₂—B grains. Upon lithiation, the anatase phase experiences a volumetric increase along the wall, while the TiO₂—B grain undergoes a contraction along its c-axis, and consequently creates a small fracture at the interface, as shown by the HAADF image in FIG. 12C. Such fractures are seen repeatedly in the post-cycling sample along the anatase walls, as shown in FIG. 14.

To confirm that Li⁺ is inserted into the bronze films during cycling, fine XRD scans are performed on a (001) Ca:TiO₂—B film grown on a (100) SrTiO₃ substrate around the strongest available diffraction peak, 006, before and after charging with Li⁺ at a rate of C/10. The results are shown in FIG. 13A. Using the SrTiO₃ substrate peaks as reference, the 006 peak of the lithiated film has shifted to lower 2θ angles, indicating a lattice expansion in the out-of-plane direction from Ca:TiO₂—B=17.98 Å to 18.04 Å. By fitting the experimental data, it is also clear that the peak has broadened as the inserted Li⁺ disturbs the crystallinity of the lattice structure. The broadening is mainly caused by local straining within the thin film as the unit cell undergoes an asymmetric deformation upon lithiation. Structural defects such as dislocations and stacking faults may also contribute to straining effects in close regions and thus to the peak broadening.

The same experiments are performed around the TiO₂—B 001 peak of a (001) TiO₂—B/Ca:TiO₂—B dual layer film grown on a (100) SrTiO₃ substrate. FIG. 13B shows that C_(TiO2-B) exhibits a slight contraction of ˜0.21% upon lithiation, which is in good agreement with the literature, where neutron diffraction results revealed a contraction in C_(TiO2-B) of 0.18%-0.49%, depending on the Li content. An obvious peak broadening is again observed.

It should be noted that the lattice constant changes observed with this procedure may not correspond to fully lithiated films. The thin films have a large surface exposed to the electrolyte and a small mass, so an unknown amount of Li charged into the film could be lost to the electrolyte via self-discharge before the cell is disassembled and the electrode is examined. Although this might also happen in the more typical powder samples of TiO₂—B, the much larger amount of active material used there may ensure that more of the Li is retained in the sample. Such a difference in sample geometry may help explain the discrepancy between our observation and the reported values in the literature. Water and anatase impurities may also have an effect on the values reported for TiO₂—B in the literature.

As discussed with regard to FIG. 12C, the larger volumetric changes of the anatase walls in the TiO₂—B films can fracture the structure during Li⁺ insertion and extraction. FIG. 14 shows a wider area of the same film, where such fracturing always occurs along the inclined anatase walls throughout the film, some at the interface of TiO₂—B and Ca:TiO₂—B and others inside the TiO₂—B phase. Therefore, this fracturing should be an actual effect of the Li⁺ insertion and extraction rather than an incidental event. A longer fracture could form between two parallel anatase walls. Because of the layered structure of TiO₂—B, all fractures are parallel to the a-b planes.

FIGS. 13A-13B also show that Li⁺ is being inserted into the films comes from XRD, where fine scans are performed around the Ca:TiO₂—B 006 peak and TiO₂—B 001 peak on respective (001) films before and after lithiation. Shifts in the peak positions indicate a slight increase in the c-axis lattice constant of Ca:TiO₂—B and a slight decrease in that of TiO₂—B. A broadening is also observed in both peaks as the inserted Li⁺ causes local straining.

As known in the art, Li plating may occur at the interface of electrolyte and Cu current collector, which would result in erroneously higher capacity measured for the Li storage material in battery cycling tests. To rule out the influence of possible Li plating on the Cu grid, a Ca:TiO₂—B thin film sample with top Cu grid current collector is charged and discharged at the rate of 1000 C for 20 cycles, taken out of the cell at a half cycle when the film was fully charged with Li⁺, and examined with an FEI Quanta scanning electron microscope. The surface image and the X-ray energy dispersive spectrum from the Cu wire are shown in FIGS. 15A-15B. No obvious Li dendrite formation is observed either on the Cu wire or on the Ca:TiO₂—B film.

In certain variations, by aligning the material to a preferred orientation, the titania-bronze structure can safely work at extremely high rates, delivering specific power of ˜20 kW kg⁻¹ at 60 C and ˜280 kW kg⁻¹ at 12000 C. Coupling with a cathode material that sustains ultrahigh rates, such as LiFePO₄ and its modifications, a superfast-charging full-cell may be made. High crystalline quality TiO₂—B or Ca:TiO₂—B films can be used to fabricate LIBs and other devices.

Example 2

Multilayered Stacks Comprising TiO₂—B Thin Films Embedded with Electrically Conductive Nanoparticles

A thin film of Pt—TiO₂—B is generated on a layer of Ca:TiO₂—B to form a multilayered stack. An image of the multilayered stack having the Pt—TiO₂—B and Ca:TiO₂—B layers is shown in FIG. 16. This image shows a substantially uniform density of Pt nanoparticles embedded within TiO₂—B. The multilayered stack has a particle concentration n of 2.82×10¹⁹ cm⁻³, a resistivity ρ of 0.033 Ω·cm, and an electron mobility μ of 6.65 cm²/Vs.

FIGS. 17A-17C show images of a multilayered stack having a SrTiO₃ substrate, a layer of Ca:TiO₂—B positioned on the substrate, and a layer of Pt—TiO₂—B positioned on the Ca:TiO₂—B layer. The Pt nanoparticles increase in relative concentration/density from a first surface of the Pt—TiO₂—B layer adjacent to the Ca:TiO₂—B layer to a second opposing surface. Rutherford Backscattered Spectroscopy (RBS) results show that the Pt—TiO₂—B layer includes about 1.5% Pt, 33.5% Ti, and about 65% O.

In FIG. 18, the rate capability of Pt—TiO₂—B thin films grown on (001) SrTiO₃ generated under vacuum and under 16 mT O₂ with channel openings on the surface is compared with the films shown in FIG. 12A. The film grown in vacuum (labeled “x,” having a high concentration/density of Pt nanoparticles) maintains energy storage capability at extremely high rates due to enhanced electrical conductivity. Such a device is suitable for high power application, such as, for example, in electric cars. The film grown in 16 mT O₂ (labeled “▴,” low density of Pt nanoparticles) has high capacity at lower rates. Such a device is suitable for applications that emphasize large capacity, but do not require high power output, such as, for example, mobile electronics. Both Pt—TiO₂—B devices can be used for environmental applications, such as, for example, for photocatalysis.

The foregoing description of the embodiments has been provided for purposes of illustration and description. It is not intended to be exhaustive or to limit the disclosure. Individual elements or features of a particular embodiment are generally not limited to that particular embodiment, but, where applicable, are interchangeable and can be used in a selected embodiment, even if not specifically shown or described. The same may also be varied in many ways. Such variations are not to be regarded as a departure from the disclosure, and all such modifications are intended to be included within the scope of the disclosure. 

What is claimed is:
 1. A composition comprising calcium-doped titania bronze (Ca:TiO₂—B), wherein the Ca:TiO₂—B is a crystalline calcium titanate in which calcium ions are uniformly distributed throughout a titanate structure.
 2. The composition according to claim 1, wherein the Ca:TiO₂—B comprises open channels aligned to out-of-plane directions.
 3. The composition according to claim 1, wherein the Ca:TiO₂—B is free from water molecules.
 4. The composition according to claim 1, wherein the Ca:TiO₂—B has a unit cell with an orthorhombic structure with a symmetry of CMCM (63).
 5. The composition according to claim 4, wherein the Ca:TiO₂—B unit cell has lattice constants of a=12.1702 Å, b=3.8013 Å, c=17.9841 Å, and α=β=γ=90°.
 6. A structure for an electronic device comprising a substrate and a layer of calcium-doped bronze (Ca:TiO₂—B), wherein the Ca:TiO₂—B is a crystalline calcium titanate in which calcium ions are uniformly distributed throughout a titanate structure.
 7. The structure according to claim 6, wherein the substrate comprises a perovskite material comprising molecules having a formula ABX₃, wherein A and B are cations of different sizes and X is an anion that bonds to both cations.
 8. The structure according to claim 7, wherein the perovskite material is selected from the group consisting of: SrTiO₃, BaTiO₃, MgSiO₃, CaTiO₃, FeTiO₃, LaMnO₃, PbTiO₃, and mixtures thereof.
 9. The structure according to claim 8, wherein the perovskite material comprises strontium titanate (SrTiO₃).
 10. The structure according to claim 7, wherein the perovskite material is doped.
 11. The structure according to claim 10, wherein the perovskite material comprises SrTiO₃ doped with niobium (Nb:SrTiO₃).
 12. The structure according to claim 6, wherein the substrate comprises a perovskite material and the layer of Ca:TiO₂—B is in direct contact with the substrate.
 13. The structure according to claim 6, wherein the substrate comprises silicon (Si).
 14. The structure according to claim 13, further comprising a buffer layer comprising a perovskite material positioned directly between the substrate and the layer of Ca:TiO₂—B.
 15. The structure according to claim 14, wherein the perovskite material is a perovskite oxide selected from the group consisting of: SrTiO₃, BaTiO₃, MgSiO₃, CaTiO₃, FeTiO₃, LaMnO₃, PbTiO₃, and mixtures thereof.
 16. The structure according to claim 6, wherein the layer of Ca:TiO₂—B is a template layer for an additional layer.
 17. The structure according to claim 16, further comprising a second layer of pure bronze (TiO₂—B) deposited on the layer of Ca:TiO₂—B.
 18. The structure according to claim 16, further comprising a second layer of titanium bronze comprising with platinum (Pt—TiO₂—B) deposited on the layer of Ca:TiO₂—B.
 19. A multilayered structure comprising: a layer of perovskite material; a layer of calcium-doped bronze (Ca:TiO₂—B) in contact with the layer of perovskite material; and a layer of pure bronze (TiO₂—B) in contact with the layer of Ca:TiO₂—B, wherein the layer of Ca:TiO₂—B is between the layer of perovskite material and the layer of TiO₂—B.
 20. The multilayered structure according to claim 19, wherein the layer of perovskite material is a substrate.
 21. The multilayered structure according to claim 20, wherein the perovskite material comprises strontium titanate (SrTiO₃) doped with niobium (Nb:SrTiO₃).
 22. The multilayered structure according to claim 19, wherein there is a near-perfect lattice match between surface structures of the Ca:TiO₂—B layer and the TiO₂—B layer.
 23. The multilayered structure according to claim 19, wherein the layer of perovskite material is directly layered onto a silicon (Si) substrate.
 24. The multilayered structure according to claim 19, wherein the layer of pure bronze (TiO₂—B) further comprises electrically conductive nanoparticles.
 25. The multilayered structure according to claim 24, wherein the electrically conductive nanoparticles are platinum nanoparticles.
 26. A lithium ion battery comprising the multilayered structure of claim
 19. 27. A method for making a multilayered structure comprising: depositing a layer of calcium-doped bronze (Ca:TiO₂—B) onto a substrate by pulsed laser deposition (PLD), wherein the PLD comprises laser ablating of a calcium titanium oxide (CaTi₄O₉) target to generate the Ca:TiO₂—B, wherein the Ca:TiO₂—B is a crystalline calcium titanate in which calcium ions are uniformly distributed throughout a titanate structure.
 28. The method according to claim 27, wherein the substrate comprises a perovskite material selected from the group consisting of: SrTiO₃, BaTiO₃, MgSiO₃, CaTiO₃, FeTiO₃, LaMnO₃, PbTiO₃, and mixtures thereof.
 29. The method according to claim 27, wherein the substrate comprises SrTiO₃ doped with niobium (Nb:SrTiO₃).
 30. The method according to claim 27, wherein the substrate comprises silicon (Si), and the method further comprises modifying the substrate before the depositing the layer of calcium-doped bronze (Ca:TiO₂—B) by first depositing a second layer of a perovskite oxide onto the substrate by epitaxial growth, wherein the perovskite oxide is selected from the group consisting of: SrTiO₃, BaTiO₃, MgSiO₃, CaTiO₃, FeTiO₃, LaMnO₃, PbTiO₃, and mixtures thereof.
 31. The method according to claim 27, further comprising forming the CaTi₄O₉ target by: i. preparing a mixture comprising greater than or equal to about 50% by weight to less than or equal to about 90% by weight TiO₂ powder with greater than or equal to 10% to less than or equal to about 50% by weight of a powder selected from CaO powder, CaTiO₃ powder, or combinations thereof; ii. sintering the mixture at a temperature greater than or equal to about 1000° C. to less than or equal to about 2000° C. to generate a powdered material; and iii. pressing the powdered material into a pellet under greater than or equal to about 5000 lbs of force to less than or equal to about 15,000 lbs of force to generate the CaTi₄O₉ target.
 32. The method according to claim 31, wherein the mixture comprises 80% by weight TiO₂ and 20% by weight CaO; the sintering is performed at a temperature of about 1400° C.; and the pressing is performed under about 10,000 lbs of force.
 33. The method according to claim 27, wherein the depositing the layer of calcium-doped bronze (Ca:TiO₂—B) by PLD comprises depositing the layer of Ca:TiO₂—B at a temperature of about 800° C.
 34. The method according to claim 27, wherein the laser ablating is performed with a 248 nm KrF excimer laser with a pulse duration of 22 ns, a fluence of about 3.4 J cm⁻², and a 10 Hz repetition rate.
 35. The method according to claim 27, further comprising depositing a layer of pure bronze (TiO₂—B) onto the layer of Ca:TiO₂—B.
 36. The method according to claim 35, wherein the depositing the layer of TiO₂—B is performed by PLD, wherein the PLD comprises laser ablating a pure TiO₂ target to generate the TiO₂—B. 